DOI: 10.1002/( adma.201604758R1) Article type: Communication

Dual-Source Precursor Approach for Highly Efficient Inverted Planar Heterojunction Perovskite Solar Cells

Deying Luo, Lichen Zhao, Jiang Wu, Qin Hu, Yifei Zhang, Zhaojian Xu, Yi Liu, Tanghao Liu, Ke Chen, Wenqiang Yang, Wei Zhang, Rui Zhu* and Qihuang Gong

D. Luo, L. Zhao, J. Wu, Q. Hu, Y. Zhang, Z. Xu, Y. Liu, T. Liu, K. Chen, W. Yang, Prof. R. Zhu, Prof. Q. GongState Key Laboratory for Artificial Microstructure and Mesoscopic Physics, Department of Physics, Peking University, Beijing, 100871, ChinaQ. Hu, W. Yang, Prof. R. Zhu, Prof. Q. GongCollaborative Innovation Center of Quantum Matter, Beijing, 100871, ChinaQ. HuMaterials Sciences Division, Lawrence Berkeley National Laboratory, CA, 94720, USADr. W. ZhangSchool of Chemistry, Joseph Banks Laboratories, University of Lincoln, Beevor Street, Lincoln LN6 7DL, United KingdomAdvanced Technology Institute, University of Surrey, Guildford GU2 7XH, United KingdomProf. R. Zhu, Prof. Q. GongCollaborative Innovation Center of Extreme Optics, Shanxi University, Taiyuan, Shanxi, 030006, China

*Corresponding Author: iamzhurui@pku.edu.cn

Keywords: Dual-Source Precursor Approach, Perovskite Solar Cells, Inverted Planar Heterojunction, Highly Efficient, Stable Mixed-Cation Perovskite

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Abstract: The highest efficiencies reported for perovskite solar cells so far have been obtained

mainly with methylammonium and formamidinium mixed cations. Currently, high-quality

mixed-cation perovskite thin films are normally made by use of anti-solvent protocols. However,

the widely used “anti-solvent”-assisted fabrication route suffers from challenges such as poor

device reproducibility, toxic and hazardous organic solvent and incompatibility with scalable

fabrication process. Here, a simple dual-source precursor approach was developed to fabricate

high-quality and mirror-like mixed-cation perovskite thin films without involving additional anti-

solvent process. By integrating the perovskite films into the planar heterojunction solar cells, a

power conversion efficiency of 20.15% was achieved with negligible current density-voltage

hysteresis. A stabilized power output approaching 20% was obtained at the maximum power

point. These results shed light on fabricating highly efficient perovskite solar cells via a simple

process, and pave the way to solar cell fabrication via scalable methods in the near future.

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Metal halide perovskites-based solar cells have rapidly risen to the forefront of the

emerging photovoltaic technologies[1-5], and the device has delivered a power conversion

efficiency (PCE) of 22.1%[6] in less than seven years of development because of their

extraordinary and unique optoelectronic properties[7-11]. The state-of-the-art perovskite solar cells

(PSCs) are normally composed of multiple functional layers[12-17] including p- and n-type charge

extraction layers, a perovskite photoactive layer and the corresponding electrodes with different

work functions. Metal halide perovskites with mixed organic cations (i.e., methylammonium

(MA) and formamidinium (FA)) have recently attracted considerable attention owing to their

lower bandgaps and higher device efficiencies than those of the single-MA-cation counterparts[18-

20]. The mixed-cation perovskite containing formamidinium, however, is more difficult to be

formed with high quality and stable phase than the methylammonium lead tri-halide (MAPbI3)

perovskites[21, 22]. This is because the α-phase FAPbI3 formed at high temperature tends to

transform back to the non-perovskite phase as the temperature decreases to room temperature[23-

25]. To address these issues, many approaches have been proposed to enhance the structure

stability of the mixed-cation perovskite films, such as cesium-induced structural transition [26, 27],

compositional engineering[18, 21, 28, 29] and additive-modulated evolution[30]. Moreover, numerous

fabrication methods including organic-cation displacement[31], anti-solvent engineering[32] as well

as intramolecular exchange[33] have been developed to yield the desired morphologies of the

mixed-cation perovskites. Among these methods, the anti-solvent engineering is representative

for “one-step” solution process that yields smooth, large-grain and pure α-phase thin films for

highly efficient mixed-cation PSCs[26, 32]. Chlorobenzene (CB) and toluene are the most

commonly used anti-solvents to precipitate perovskite or its intermediate phase from the

perovskite solution containing mixing solvents[23,29], e.g., dimethylsulfoxide (DMSO), γ-

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butyrolactone (GBL) and N,N-dimethylformamide (DMF)[26, 32]. Although the use of organic anti-

solvents is helpful to fabricate high-quality mixed-cation perovskite thin films, the toxicity of the

anti-solvents would hinder their industrial application[34]. Meanwhile, the anti-solvent is usually

dropped on the center of the substrate which leads to spatially inhomogeneous nucleation on the

whole perovskite films, unfavorable for large-area solar cell fabrication and device

reproducibility[35]. The anti-solvent bath has provided a potential solution to address the

challenge of spatially inhomogeneous nucleation at the surface of the thin film [16]. Furthermore,

the anti-solvent-assisted fabrication protocol is not compatible with the existing scalable

methods, e.g., slot-die coating and roll-to-roll fabrication.

As a result, it is warranted to develop a simple and effective method to fabricate high-

quality mixed-cation perovskite thin films for the application in the solar cells. Huang et al. have

demonstrated that pure α-phase mixed-cation perovskites can be obtained on the preheated

substrates (> 140 °C) by doctor balding, achieving a PCE of 18%[22]. Zhao et al. have proposed a

hydrochloric acid-accelerated formation of planar perovskite thin film for the PSCs without anti-

solvent[36], with a PCE less than 15%. Recently, Grätzel et al. have developed a vacuum flash-

assisted solution process by the removal of regular anti-solvent used in the “one-step” solution

process, and ultimately boosted a certified PCE to 19.6% on the large-scale substrate (>1 cm2)[34].

The improved photovoltaic performance critically depends on large-grain and compact films

with excellent optoelectronic properties. Recently, Zhang et al. have reported a MAI-based

“acetate route” by replacing the regular lead halides (PbCl2, PbBr2 or PbI2) with lead acetate

(PbAc2) as the lead source, enabling smooth MAPbI3 thin films by “one-step” solution process

with much shorter annealing time[37, 38]. Afterwards, we have further demonstrated that the

incorporation of bromide (MABr) as additive in the “acetate route” enabled to obtain

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controllable ultra-smooth and large-grain films without any additional anti-solvents[39], and a

reproducible device efficiency exceeding 18% in the inverted planar heterojunction PSCs was

achieved. However, this simple approach was not readily applicable to the case of mixed-cation

perovskites (MA+/FA+-based), since the undesired non-perovskite phase always appeared after

thermal annealing at 100 – 170 °C.

To overcome these challenges, herein we develop a simple and convenient dual-source

precursor approach (DSPA) to produce high-quality perovskite films with mixed MA+/FA+

cations, where the dual sources of PbAc2:MAI (1:3) and PbI2:FAI (1:1) are dissolved into DMF

solvent. The existing MAI-based “acetate route” within the solution makes sure that DMF

solvent couldn’t damage the perovskite films during the film coating process. Synchronous

evaporation of MA·Ac and DMF at the elevated-temperature duration induces rapid

crystallization and yields pure α-phase perovskite at an appropriate temperature. The absence of

pinholes within the pure α-phase perovskite films enables superior bulk charge transport. As a

demonstration, by integrating uniform and mirror-like smooth perovskite films into the planar

heterojunction PSCs, a power conversion efficiency exceeding 20%, was achieved with

negligible hysteresis effect. The device also showed a stabilized power output approaching 20%

at the maximum power point (MPP). Over 80% of the initial PCE for the champion solar cell can

be maintained after the device was placed in the dark condition for 200 h.

Figure 1a shows the ultraviolet–visible (UV–vis) absorption spectra of the thin films

prepared with different perovskite precursor solutions. Compared with the single-cation MAPbI3

films formed by the “acetate route”, the perovskite films obtained with the dual-source precursor

approach had a broad absorption in the UV–vis light wavelength range. The increased absorption

edge was consistent with the reference film based on PbI2:MAI:FAI (1:0.6:0.4). For the films

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made by the PbAc2:FAI (1:3) and PbAc2:MAI:FAI (1:1.8:1.2) precursor solutions, the short

wavelength absorption indicated that non-perovskite phases were formed with these two recipes.

The further information on the solutions and the deposited-films are shown in Figure S1

(supporting information). To further understand the dual-source solution, the measured FTIR

spectra show that no new functional groups appeared in the solution with PbAc2:PbI2:MAI:FAI

(Figure S2, supporting information), which means that the MAI-based “acetate route” could be

sustained in the dual-source precursor solution. X-ray diffraction (XRD) analysis in Figure 1b

confirmed that the films with non-perovskite phase were primarily composed of δ-FAPbI3 and

PbI2-complex intermediate, because the FA-based non-perovskite can be easily formed at room

temperature[23-25]. Furthermore, XRD illustrates that the mixed-cation MA0.60FA0.40PbI3 perovskite

obtained from the DSPA procedure at a low temperature (i.e., 120 °C) had purer perovskite

phase (Figure 1b). Its diffraction peaks shifted to a smaller angle relative to pristine MAPbI3.

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Figure 1. a) The ultraviolet–visible (UV–vis) absorption spectra and b) X-ray diffraction (XRD) patterns of the perovskite thin films synthesized from different perovskite precursor solutions through annealing at 120 °C for 10 min, where the utilized recipes were PbAc2:MAI (1:3), PbAc2:FAI (1:3), PbAc2:MAI:FAI (1:1.8:1.2), PbAc2:PbI2:MAI:FAI (0.6:0.4:1.8:0.4) and PbI2:MAI:FAI (1:0.6:0.4), respectively. c) The crystal evolution of the perovskite thin films obtained from the recipe of PbI2:MAI:FAI (1:0.6:0.4) undergoing annealing at room temperature (RT), 120 °C and 140 °C, respectively.

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To further elucidate the mechanism that governs the phase transition of the mixed-cation

perovskite MA0.60FA0.40PbI3 by using DSPA, as a control recipe, PbI2:MAI:FAI was studied as

well. The mixed-cation perovskite film with a rough surface was obtained as shown in Figure S1

(supporting information). XRD results in Figure 1c illustrate the temperature-induced structural

variation in the thin films. The non-perovskite δ-phase for mixed-cation thin film was formed at

room temperature, and the α-phase perovskite and intermediate phase co-existed in the film

formed at 120 °C. Ultimately stable and pure perovskite phase was achieved at 140 °C. These

results suggest that the phase-transition temperature from hexagonal to pseudo-cubic phase for

the mixed-cation perovskite is lower than that of the pure FAPbI3 (175 °C)[25], probably due to

the small radius of MA+ enabling to tune the tolerance factor[23] approaching to stable cubic

region. By using the DSPA approach, the phase-transition temperature could decrease to 120 °C

and the ultra-smooth morphological features based on the “acetate route”could be sustained.

To get a physical insight into the quality of mixed-cation perovskites via the DSPA

process, the impact of crystal evolution upon thermal annealing temperatures has been explored

as well. As shown in Figure 2a, mixed phases of non-perovskite (δ-phase) and perovskite (α-

phase) were formed after annealing at room temperature with the residual DMF vapor. At the

annealing temperature of 100 °C, mixed phases composed of tetragonal structure (MAPbI3)[40]

and δ-phase appeared and the major diffraction peaks were indexed to MAPbI3. At the annealing

temperature of 120 °C, diffraction peaks attributed to the δ-phase totally disappeared and perfect

reflections from (110) and (220) corresponding to the α-phase appeared, without obvious change

at a higher annealing temperature (140 °C).

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Figure 2. a) Comparisons of the MA0.60FA0.40PbI3 obtained via the dual-source precursor approach (DSPA) undergoing different thermal annealing temperatures. b) Optical absorption spectra of (MA)1-x(FA)xPbI3 (x =0, 0.25, 0.30, 0.33, 0.40) annealed at 140 °C. c) Illustration of thin film growth mechanism for the mixed-cation perovskite via the DSPA method.

Based on these results, we propose a thin film growth mechanism as shown in Figure 2c.

After spin-coating, five species including PbI64-, MA+, FA+, Ac- and residual DMF co-existed in

the as-deposited films. With the aid of residual DMF vapor, intercalation of FA+ into PbI64- (from

the adducts of PbI64-·DMF) would lead to the formation of FAPbI3 at room temperature which

was confirmed by the XRD characterization. The as-deposited films were instantly heated on a

hot plate, then the residual DMF and MA·Ac could be driven out simultaneously, producing

MAPbI3 and FAPbI3 synchronously at an elevated temperature. The nucleation of α-phase

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FAPbI3 needs to be at a higher temperature[25], thus the pure α-phase only can be obtained at an

appropriate temperature (e.g., > 120 °C). Ultimately, the mixed-cation perovskite films exhibited

single α-phase owing to the intercalation of the MA+ and FA+ cations into the same crystal unit.

The major diffraction peaks had a gradual shift to smaller diffraction angle than those of the

pristine MAPbI3 accompanying with the increasing of FA+ fraction, as the radius of FA cation is

larger than MA cation (Figure S3, supporting information).

To deeply understand the variations in optical properties after the FA cations were

introduced into the pristine MAPbI3, we determined the optical absorption of (MA)1-x(FA)xPbI3

mixed-cation perovskite where x was 0, 0.25, 0.30, 0.33 and 0.40, respectively, as shown in

Figure 2b. As expected, the optical absorption onset wavelength of the perovskite film could be

tuned from 780 nm for the pristine MAPbI3 to 815 nm for the MA0.60FA0.40PbI3, confirming the

incorporation of FA cations into the perovskite crystals. These observations were in agreement

with XRD analysis before. Besides of crystal structure and optical absorption for the mixed-

cation perovskite, the surface morphologies of the perovskite thin films are widely recognized as

a key factor to govern ultimate device performance. Ultra-smooth and compact perovskite films

are truly important to achieve highly efficient PSCs. Scanning electron microscopy (SEM)

images of the mixed-cation perovskite MA0.60FA0.40PbI3 thin films were measured after thermal

annealing at 100 °C, 120 °C and 140 °C for 10 min, respectively. Figure 3a shows that bright

crystals or islands randomly existed on the top of the perovskite films obtained at the annealing

temperature of 100 °C, indicating the phase segregation as evidenced by XRD characterization.

The phase segregation was more obvious after RT annealing, as well as for the case of

PbAc2:MAI:FAI without PbI2 (Figure S4, supporting information). After annealing at 120 °C,

however, phase segregation was basically eliminated, resulting in uniform and compact

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perovskite thin films with large crystal grains. After even higher annealing temperature (140 °C),

the mixed-cation perovskite thin films will not suffer from disintegration (Figure 2a), and even

smoother films were obtained across the whole substrate surface.

Figure3. a) Top-view SEM images of MA0.60FA0.40PbI3 perovskite after annealing at 100 °C, 120 °C and 140 °C for 10 min, respectively, with the scale bar of 500 nm. b) Conductive atomic force microscopy (c-AFM) characterizations in contact mode for the α-phase perovskite films with different MA+/FA+ ratios: 3:1, 2:1 and 3:2. Given the small values of surface current, the mapping of current value vs. color is built by evaluating the logarithm of current value.

Recently, it has been reported that charge-carrier mobility and diffusion length for

MAPbI3 and FAPbI3 perovskites are not identical[27, 41]. Thus, it is necessary to explore its

conductivity, mobility and diffusion length for the mixed-cation perovskite thin film. At first,

conductive atomic force microscopy (c-AFM) characterizations in contact mode for the α-phase

perovskite films with the varied mixing ratios (MA+/FA+) were performed to investigate the

distribution of the conductivity, as shown in Figure 3b. The surface AFM images of the

perovskites with different MA+/FA+ ratios annealed at 140 °C are shown in Figure S5, supporting

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information. In spite of the uniform conductivity of all the mixed-cation perovskites, as revealed

by c-AFM characterization, the relatively high surface current were obtained at the optimized

ratio of MA+/FA+ at 3:2. The high surface current could be attributed to large grains and pure

phase for the mixed-cation perovskite thin film. These suggest that the mixed-cation perovskite

thin film with the MA+/FA+ at 3:2 has a good electrical conductivity.

To mimic the recombination of photo-generated carriers during the PSCs device

operation, we presented time-resolved photoluminescence (TRPL) decays, measuring the peak

emission at 770 nm. The measured results for the mixed-cation perovskite with pure α-phase and

mixed phases are shown in Figure 4a, b. This allowed us to derive information on the dynamics

of charge carrier transport in the perovskite films or extraction at the interfaces. The measured

results were fitted with a bi-exponential model in the presence or absence of the quenching

layers, suggesting that bi-molecular recombination mechanism of photo-generated charge

carriers is predominant for the mixed-cation system. Comparing the mixed phases (α, δ) with

pure α-phase of MA0.60FA0.40PbI3, the pure α-phase MA0.60FA0.40PbI3 perovskite had a long

lifetime, exhibiting a time constant (τ) 231 ± 7 ns. The average diffusion length LD of the species

was then further determined from a classical diffusion model[9]. An electron-carrier diffusion

length of 627 ± 9 nm and a hole-carrier diffusion length of 504 ± 8 nm were achieved in the

identical film (See details in Table. S1). These diffusion lengths were comparable to the pristine

MAPbI3 resulting from PbAc2-based system[42]. By contrast, the presence of non-perovskite (δ-

phase) within the mixed-cation perovskite thin film, even at a very small fraction, led to halved

charge carrier lifetime and diffusion length as shown in Figure 4a. Moreover, the addition of the

phenyl-C61-butyric acid methyl ester (PC61BM) and poly[bis(4-phenyl)(2,4,6-

trimethylphenyl)amine] (PTAA) layers for pure α-phase perovskite accelerated the TRPL decay,

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with the observed τ of 2.69 ± 0.1 ns and 5.9 ± 0.1 ns, respectively. The fast and balanced charge-

carrier extraction rates suggest that PTAA and PC61BM can serve as excellent charge-carrier

extraction layers at the interfaces, owing to their favorable energy level alignments with the

perovskite semiconductor (Figure 4c). These results are consistent with the previously published

results[43, 44].

Current experimental observations suggest that there might be room to improve these

remarkable device efficiencies via the optimization of charge-carrier mobility. The charge-carrier

mobility was extracted from current-electric field characteristics of the single-carrier devices.

This method was commonly used to derive the information on mobility in semiconductors[45-47].

Current values vs. electric fields from the two single-carrier devices are shown in Figure 4d on a

logarithmic scale. The architecture of the electron-only single-carrier device was

ITO/TiO2/MA0.60FA0.40PbI3/PC61BM/BCP/Ag (electron-only device), while the hole-only single-

carrier device was composed of ITO/PTAA/MA0.60FA0.40PbI3/PTAA/Ag (hole-only device). As

we mentioned above, the MA0.60FA0.40PbI3 layer was capable of forming favorable energy-level

alignments (i.e., Ohmic contact) with PTAA and PC61BM quenchers which ensure only electron

or hole carriers left in the perovskite. According to the current-electric field (CF) curves, we

found the CF curves could be classified as three different regions, e.g., linear Ohmic regime,

trap-filled regime and space charge limited current (SCLC) regime. The SCLC was regarded as

the maximum current that can be sustained within semiconductor materials. An Ohmic contact is

therefore an ideal interface, which is capable of injecting enough charges from the electrodes to

the semiconducting materials to sustain SCLC especially when trap states are negligible. In this

case, the SCLC for unipolar transport within the single-carrier devices can be given by a Mott-

Gurney law[45, 47]. Estimating mobility via a linear fit based on the Mott-Gurney law led to the

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estimated values of 18.38 cm2/(V·s) (electron) and 19.96 cm2/(V·s) (hole), respectively, and the

order of the magnitudes were consistent with other results[48]. The balanced charge-carrier

mobility provides balanced transport within the perovskite layer, which will result in a high fill

factor for the PSCs.

Figure 4. a) Time-resolved photoluminescence (TRPL) spectra of the MA0.60FA0.40PbI3

perovskite films of the mixed (α, δ) phases and α-phase without or b) with electrons (PC61BM)

and holes (PTAA) quenchers at 770 nm, solid lines are the fitted curves by a bi-exponential

model. c) Schematic illustration of charge transport and charge transfer. d) Current-electric field

curves of the single-carrier devices, ITO (indium tin oxide)/TiO2/MA0.60FA0.40PbI3/PC61BM/Ag

(electron-only device); ITO/PTAA/MA0.60FA0.40PbI3/PTAA/Ag (hole-only device).

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As shown in Figure 5a, the inverted planar heterojunction PSCs in this study were

adapted as ITO (indium tin oxide)/doped-PTAA (10 nm)/MA1-xFAxPbI3 (300~330 nm, x = 0.25,

0.30, 0.33 and 0.40)/PC61BM (30 nm)/C60 (20 nm)/BCP (7 nm)/Ag (80 nm), where 2,9-dimethyl-

4,7- diphenyl-1,10-phenanthroline (BCP) served as a hole-blocking layer. A bilayer of

PC61BM/C60 was used as the electron extraction layer while PTAA was the hole extraction layer

with a matching energy level and superior optical transmission. In order to ascertain the optimum

ratio of MA+ to FA+ in the mixed-cation perovskite, we fabricated a series of inverted-

architecture PSCs composed of perovskite films with varying MA+/FA+ ratios. Table S2

summarizes the device performances measured under the simulated AM 1.5G 100 mW/cm2

illumination. As the fraction of FA+ increased from 0.25 to 0.40, the short-circuit photocurrent

density (Jsc) was continuously improved without sacrificing open-circuit voltage (Voc) and fill

factor (FF). The optimal device performance was achieved when the ratio of MA+/FA+ was 3:2.

Figure 5b shows the cross-sectional SEM image of the perovskite layer deposited on the

ITO/doped-PTAA substrate, which was covered with PC61BM/C60/BCP multiple layers. The

MA0.60FA0.40PbI3 perovskite thin film made by the DSPA was uniform and the grain size in the

vertical direction was comparable to the perovskite film thickness. This indicates that the photo-

generated charge carrier can efficiently transport across the mixed-cation perovskite films and

reach the corresponding charge carrier extraction interfaces before recombination. Moreover,

tens of nanometers of PC61BM layer can totally cover the whole perovskite photoactive layer,

which means the photo-generated carriers from the perovskite layer can be effectively extracted

by the PC61BM interfacial layer.

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Figure 5. a) Schematic diagram of an inverted planar heterojunction perovskite solar cell. b) Cross-sectional SEM image of the perovskite film fabricated by the DSPA approach, and the scale bar is 280 nm. c) The best device performance. d) Incident photon-to-electron-conversion efficiency (IPCE) spectrum of the ITO/doped-PTAA/MA0.60FA0.40PbI3/PC61BM/C60/BCP/Ag under simulated AM 1.5G illumination of 100 mW/cm2. e) Current density-voltage (J-V) curves of the device measured with forward and reverse scans under the illumination. f) The stabilized power output at the maximum power point (MPP).

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Based on the conventional inverted-architecture PSCs integrating with the optimized

perovskite layers, a champion PCE of 20.15% was achieved after optimization, with a Jsc of

23.11 mA/cm2, a Voc of 1.09 V and a high FF of 0.80, as shown in Figure 5c. The incident

photon-to-electron conversion efficiency (IPCE) spectrum in Figure 5d demonstrated the high

quantum conversion efficiency throughout the entire wavelength range. The integrated

photocurrent (22.10 mA/cm2) was in agreement with the Jsc from the J-V scanning, and the

optimized device had a negligible photocurrent hysteresis (Figure 5e) by eliminating the trap

states at the grain boundaries through the fullerene passivation effects. The improved perovskite

films synthesized with the DSPA approach had long diffusion lengths and high charge-carrier

motilities, which facilitated the extraction of charge carriers at the corresponding interfaces

before recombination. All of these improvements could contribute to high IPCE value and device

efficiency. Ultimately, the steady-state measurements for Jsc and PCE extracted at the MPP are

shown in Figure 5f, which confirmed the accuracy of device parameters extracted from the J-V

curves. To demonstrate the reproducibility of device performances for PSCs fabricated through

the DSPA route, we fabricated 40 devices from different batches and presented the histograms of

their average values in Figure S6 (supporting information). Most of devices have achieved PCEs

between 19% and 20%, among which quite a few were over 20%, suggesting the improved

reproducibility of this DSPA method. Considering the potential instability issues induced by the

n-type PC61BM interface, the phase stability of the optimized PSCs without encapsulation was

studied in the nitrogen atmosphere. Over 80% of the initial PCE for the champion device could

be maintained after the device was kept in the dark condition for 200 h (Figure S6, supporting

information). In contrast, the PCE of the control device made by the regular one-step method

degraded to < 50% of its initial value within 100 h with storage in the nitrogen atmosphere. This

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suggests that the stability of the mixed-cation perovskite photoactive layer has been improved by

the proposed method in this work.

In summary, a simple and effective dual-source precursor approach was developed to

fabricate mixed-cation perovskite thin film with pure and stable crystal phase at a relatively low

annealing temperature, yielding uniform and compact perovskite thin films with excellent

optoelectronic properties. Ultimately, a champion PCE of 20.15% was achieved in the inverted

planar heterojunction perovskite solar cells based on the DSPA fabrication method, and a

stabilized output efficiency of 19.80% was obtained at the maximum power point. The optimized

device showed negligible J-V hysteresis. These results suggest that the DSPA is a promising

method to fabricate the high-performance perovskite solar cells. Moreover, the DSPA process

simplifies the device fabrication steps and avoids the use of toxic anti-solvent, thus opens a new

avenue to the scalable fabrication of PSCs in the near future.

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Experimental Section

Materials and solution preparation: Methylamine iodide (MAI) and formamidinium

iodide (FAI) were synthesized using the methods reported by the previous literatures[3, 18]. Lead

acetate trihydrate (PbAc2·3H2O, CAS no. 6080-56-4) was purchased from Sinopharm Group

Company and dehydrated before use. Poly[bis(4-phenyl)(2,4,6-trimethylphenyl)amine] (PTAA)

was purchased from Xi’an Polymer Light Technology Corp. (China), and [6,6]-phenyl-C61-

butyric acid methyl ester (PC61BM) was purchased from C-nano Tech. (USA). The 2,9-dimethyl-

4,7-diphenyl-1,10-phenanthroline (BCP) and 2,3,5,6-Tetrafluoro-7,7,8,8-

tetracyanoquinodimethane (F4-TCNQ) were purchased from Jilin OLED company (China).

Besides, all the liquid reagents including N,N-dimethylformamide (DMF) and chlorobenzene

(CB) were purchased from commercial sources (Acros) and used as received. The dual-source

solution (1 M Pb2+ (PbAc2 and PbI2) in DMF) was prepared in a glovebox by dissolving the two

(PbAc2:MAI (1:3) and PbI2:FAI (1:1) in DMF, respectively) precursor solutions together. It was

slowly stirred for 30 min until a transparent yellow solution yielded. The mole ratio of

PbAc2:PbI2:MAI:FAI was 0.6:0.4:1.8:0.4. Details about the preparation of dual-source precursor

solution are given in the supporting information. For comparisons, PbAc2:FAI (1:3),

PbAc2:MAI:FAI (1:1.8:1.2), PbAc2:MAI (1:3) and PbI2:MAI:FAI (1:0.6:0.4) were dissolved into

DMF with fixed mole ratios, respectively. F4-TCNQ-doped PTAA solution was prepared with

the reported methods[49, 50], and the total concentration of solids was around 0.5 wt% relative to

CB solvent. Similarly, PC61BM solution in CB solvent was prepared by dissolving 20 mg

PC61BM into CB solvent and stirred using a magnet over 2 h at 70 °C.

Device fabrication: The inverted-architecture planar heterojunction PSCs were fabricated

on the pre-patterned ITO substrates. The pre-patterned ITO substrates were ultrasonically

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cleaned with diluted detergent, deionized water, acetone, and isopropyl alcohol (IPA) in

succession. The as-cleaned substrates were heated in an oven at 60 °C for overnight followed by

UV-Ozone treatment for 10 min. 0.5 wt% F4-TCNQ-doped PTAA solution was spin-coated onto

the ITO substrates at 4000 rpm for 30 s, and the samples were then annealed at 120 °C for 20

min in a N2-filled glovebox with H2O and O2 concentrations of < 0.1 ppm (at room temperature).

The perovskite precursor solution was spin-coated on the top of PTAA-modified ITO substrate at

4000 rpm for 25 s. The samples were further annealed on a hot plate at 100 °C, 120 °C or 140 °C

for 10 min. Next, 35 μL of PC61BM (20 mg/mL in CB) solution was then spin-coated on the top

of perovskite layer at 1000 rpm for 30 s to form an electron extracting layer. After that, the

samples were transferred to a vacuum chamber without exposing to air. Finally, the rest of layer-

by-layer architectures composed of C60 (20 nm), BCP (7 nm) were thermally evaporated in a

vacuum chamber with the base pressure of < 4×10-4 Pa. The metal silver electrode (80 nm) was

evaporated in the other vacuum chamber (< 4×10-4 Pa) through a metal mask of 0.09 cm2. For the

single charge-carrier devices, the fabrication procedures were distinct. Electron-only device: 20

nm-thick TiO2 layer was spin-coated on the top of ITO substrate from a dispersion solution of

TiO2 nanoparticles (5 mg/mL) followed by heating at 130 °C for 5 min. The perovskite and

PC61BM layers together with metal electrode were the same as the fabrication of solar cells.

Hole-only device: First, the PTAA layer and the mixed-cation perovskite layer were spin-coated

on the ITO substrates in succession, where the procedures were the same as the fabrication of

solar cells. Next, the 0.5 wt% F4-TCNQ-doped PTAA solution was spin-coated on the top of the

as-prepared mixed-cation perovskite layer at 4000 rpm for 30 s. After that, the silver electrode

(80 nm) was then thermally evaporated on the top of the PTAA layer at a pressure of < 4×10-4 Pa

through a metal shallow mask.

20

Device measurements: The solar cells were measured under simulated solar irradiation by

a 150 W class AAA solar simulator (XES-40S1, SAN-EI) with an AM 1.5G filter. Light

intensity of 100 mW/cm2 was calibrated by use of a standard monocrystalline silicon solar cell

with a KG-5 filter. The J-V curves for the solar cells were recorded with a Keithley2400 Source

Meter. To ensure the accuracy of device efficiency, an aperture mask with an area of 0.07 cm2

was exploited during the measurements. J-V measurements of the PSCs without encapsulation

were carried out in a N2-filled glovebox. The measuring condition was: reverse scan (1.2 V → -

0.1 V, scan rate 10 mV/s, and no delay time) and forward scan (-0.1 V → 1.2 V, scan rate 10

mV/s, and no delay time). The devices with encapsulation were taken out for IPCE

measurements using Zolix Instruments (China). The current-voltage characteristics for the

single-carrier devices under the darkness were measured using the electrochemical workstation

(Autolab PGSTAT302N, Metrohm, Switzerland).

Other characterizations: The ultraviolet–visible (UV–Vis) absorption spectra of mixed-

cation perovskite films were measured by a spectrophotometer (UH4150, Hitachi, Japan).

Scanning electron microscopy (SEM) images were obtained through a field-emission SEM (FEI

Nova Nano SEM 430). The X-ray diffraction (XRD) samples prepared on ITO substrates were

characterized by Mini Flex 600 (Rigaku, Japan). The atomic force microscopy (AFM) images

were collected by an atomic force microscopy from Thchcomp (5100N, Hitachi, Japan), and the

conductive atomic force microscopy (c-AFM) images were carried out in dark. The TRPL

spectra were measured at 770 nm upon excitation at 470 nm via a fluorescence

spectrophotometer (FLS980, Edinburgh Instruments, England). The film thickness was measured

by stylus profilometry (Bruker Dektak XT, Germany).

21

Supporting Information

Supporting Information is available from the Wiley Online Library or from the author.

Acknowledgements

D.L. and L.Z. contributed equally to this work. This work was financially supported by the 973

Program of China (2015CB932203), the National Natural Science Foundation of China

(61377025 and 91433203), and the Young 1000 Talents Global Recruitment Program of China.

Q.H. also received support from the Advanced Light Source Doctoral Fellowship in Residence at

the Lawrence Berkeley National Laboratory.

Received: ((will be filled in by the editorial staff))Revised: ((will be filled in by the editorial staff))

Published online: ((will be filled in by the editorial staff))

References

[1] A. Kojima, K. Teshima, Y. Shirai, T. Miyasaka, J. Am. Chem. Soc. 2009, 131, 6050.

[2] M. M. Lee, J. Teuscher, T. Miyasaka, T. N. Murakami, H. J. Snaith, Science 2012, 338,

643.

[3] J. Burschka, N. Pellet, S. J. Moon, R. Humphry-Baker, P. Gao, M. K. Nazeeruddin, M.

Grätzel, Nature 2013, 499, 316.

[4] D. Y. Son, J. W. Lee, Y. J. Choi, I. H. Jang, S. Lee, P. J. Yoo, H. Shin, N. Ahn, M. Choi,

D. Kim, N. G. Park, Nat. Energy 2016, 1, 16081.

[5] A. Mei, X. Li, L. Liu, Z. Ku, T. Liu, Y. Rong, M. Xu, M. Hu, J. Chen, Y. Yang, M.

Grätzel, H. Han, Science 2014, 345, 295.

[6] NREL chart, http://www.nrel.gov/pv/assets/images/efficiency-chart.png.

[7] L. M. Pazos-Outon, M. Szumilo, R. Lamboll, J. M. Richter, M. Crespo-Quesada, M. Abdi-

Jalebi, H. J. Beeson, M. Vrucinic, M. Alsari, H. J. Snaith, B. Ehrler, R. H. Friend, F.

Deschler, Science 2016, 351, 1430.

22

[8] J. S. Manser, P. V. Kamat, Nat. Photon. 2014, 8, 737.

[9] S. D. Stranks, G. E. Eperon, G. Grancini, C. Menelaou, M. J. P. Alcocer, T. Leijtens, L. M.

Herz, A. Petrozza, H. J. Snaith, Science 2013, 342, 341.

[10] G. Xing, N. Mathews, S. Sun, S. S. Lim, Y. M. Lam, M. Grätzel, S. Mhaisalkar, T. C.

Sum, Science 2013, 342, 344.

[11] T. Liu, K. Chen, Q. Hu, J. Wu, D. Luo, S. Jia, R. Zhu, Q. Gong, Chin. Chem. Lett. 2015,

26, 1518.

[12] W. Zhang, G. E. Eperon, H. J. Snaith, Nat. Energy 2016, 1, 16048.

[13] M. A. Green, A. Ho-Baillie, H. J. Snaith, Nat. Photon. 2014, 8, 506.

[14] T. Liu, K. Chen, Q. Hu, R. Zhu, Q. Gong, Adv. Energy Mater. 2016, 6, 1600457.

[15] Q. Hu, J. Wu, C. Jiang, T. Liu, X. Que, R. Zhu, Q. Gong, ACS Nano 2014, 8, 10161.

[16] Y. Zhou, M. Yang, W. Wu, A. L. Vasiliev, K. Zhu, N. P. Padture, J. Mater. Chem. A 2015,

3, 8178.

[17] T. Liu, Q. Hu, J. Wu, K. Chen, L. Zhao, F. Liu, C. Wang, H. Lu, S. Jia, T. Russell, R. Zhu,

Q. Gong, Adv. Energy Mater. 2016, 6, 1501890.

[18] N. J. Jeon, J. H. Noh, W. S. Yang, Y. C. Kim, S. Ryu, J. Seo, S. I. Seok, Nature 2015, 517,

476.

[19] H. H. Fang, F. Wang, S. Adjokatse, N. Zhao, J. Even, M. A. Loi, Light Sci. Appl. 2016, 5,

e16056.

[20] N. Pellet, P. Gao, G. Gregori, T. Y. Yang, M. K. Nazeeruddin, J. Maier, M. Grätzel,

Angew. Chem. Int. Ed. 2014, 53, 3151.

[21] J. Liu, Y. Shirai, X. Yang, Y. Yue, W. Chen, Y. Wu, A. Islam, L. Han, Adv. Mater. 2015,

27, 4918.

[22] Y. Deng, Q. Dong, C. Bi, Y. Yuan, J. Huang, Adv. Energy Mater. 2016, 6, 1600372.

[23] A. Binek, F. C. Hanusch, P. Docampo, T. Bein, J. Phys. Chem. Lett. 2015, 6, 1249.

[24] Y. Zhou, Z. Zhou, M. Chen, Y. Zong, J. Huang, S. Pang, N. P. Padture, J. Mater. Chem. A

2016, 4,17623.

[25] J. A. Aguiar, S. Wozny, T. G. Holesinger, T. Aoki, M. K. Patel, M. Yang, J. J. Berry, M.

Al-Jassim, W. Zhou, K. Zhu, Energy Environ. Sci. 2016, 9, 2372.

23

[26] M. Saliba, T. Matsui, J. Y. Seo, K. Domanski, J. P. CorreaBaena, M. K. Nazeeruddin, S.

M. Zakeeruddin, W. Tress, A. Abate, A. Hagfeldt, M. Grätzel, Energy Environ. Sci. 2016,

9, 1989.

[27] D. P. McMeekin, G. Sadoughi, W. Rehman, G. E. Eperon, M. Saliba, M. T. Hörantner, A.

Haghighirad, N. Sakai, L. Korte and B. Rech, M. B. Johnston, L. M. Herz, H. J. Snaith,

Science 2016, 351, 151.

[28] Y. Jo, K. S. Oh, M. Kim, K. H. Kim, H. Lee, C. W. Lee, D. S. Kim, Adv. Mater. Inter.

2016, 3, 1500768.

[29] D. Bi, W. Tress, M. I. Dar, P. Gao, J. Luo, C. Renevier, K. Schenk, A. Abate, F. Giordano

and J. P. C. Baena, J. D. Decoppet, S. M. Zakeeruddin, M. K. Nazeeruddin, M. Grätzel, A.

Hagfeldt, Sci. Adv. 2016, 2, e1501170.

[30] Z. Wang, Y. Zhou, S. Pang, Z. Xiao, J. Zhang, W. Chai, H. Xu, Z. Liu, N. P. Padture, G.

Cui, Chem. Mater., 2015, 27, 7149.

[31] Y. Zhou, M. Yang, S. Pang, K. Zhu, N. P. Padture, J. Am. Chem. Soc. 2016, 138, 5535.

[32] N. J. Jeon, J. H. Noh, Y. C. Kim, W. S. Yang, S. Ryu, S. I. Seok, Nat. Mater. 2014, 13,

897.

[33] W. S. Yang, J. H. Noh, N. J. Jeon, Y. C. Kim, S. Ryu, J. Seo, S. I. Seok, Science 2015,

348, 1234.

[34] X. Li, D. Bi, C. Yi, J. D. Décoppet, J. Luo, S. M. Zakeeruddin, A. Hagfeldt, M. Grätzel,

Science 2016, 353, 58.

[35] W. Chen, Y. Wu, Y. Yue, J. Liu, W. Zhang, X. Yang, H. Chen, E. Bi, I. Ashraful, M.

Grätzel, L. Han, Science 2015, 350, 944.

[36] G. Li, T. Zhang, Y. Zhao, J. Mater.Chem. A 2015, 3, 19674.

[37] W. Zhang, M. Saliba, D. T. Moore, S. K. Pathak, M. T. Horantner, T. Stergiopoulos, S. D.

Stranks, G. E. Eperon, J. A. Alexander-Webber, A. Abate, A. Sadhanala, S. Yao, Y. Chen,

R. H. Friend, L. A. Estroff, U. Wiesner, H. J. Snaith, Nat. Commun. 2015, 6, 6142.

[38] W. Zhang, S. Pathak, N. Sakai, T. Stergiopoulos, P. K. Nayak, N. K. Noel, A. A.

Haghighirad, V. M. Burlakov, D. W. deQuilettes, A. Sadhanala, W. Li, L. Wang, D. S.

Ginger, R. H. Friend, H. J. Snaith, Nat. Commun. 2015, 6, 10030.

[39] L. Zhao, D. Luo, J. Wu, Q. Hu, W. Zhang, K. Chen, T. Liu, Y. Liu, Y. Zhang, F. Liu, T. P.

Russell, H. J. Snaith, R. Zhu, Q. Gong, Adv. Funct. Mater. 2016, 26, 3508.

24

[40] C. Quarti, E. Mosconi, J. M. Ball, V. D'Innocenzo, C. Tao, S. Pathak, H. J. Snaith, A.

Petrozza, F. De Angelis, Energy Environ. Sci. 2016, 9, 155.

[41] Q. Han, S. H. Bae, P. Sun, Y. T. Hsieh, Y. M. Yang, Y. S. Rim, H. Zhao, Q. Chen, W. Shi,

G. Li, Y. Yang, Adv. Mater. 2016, 28, 2253.

[42] J. Xu, A. Buin, A. H. Ip, W. Li, O. Voznyy, R. Comin, M. Yuan, S. Jeon, Z. Ning, J. J.

McDowell, P. Kanjanaboos, J. P. Sun, X. Lan, L. N. Quan, D. H. Kim, I. G. Hill, P.

Maksymovych, E. H. Sargent, Nat. Commun. 2015, 6, 7081.

[43] Y. Shao, Y. Yuan, J. Huang, Nat. Energy 2016, 1, 15001.

[44] Y. Shao, Z. Xiao, C. Bi, Y. Yuan, J. Huang, Nat. Commun. 2014, 5, 5784.

[45] D. Shi, V. Adinolfi, R. Comin, M. Yuan, E. Alarousu, A. Buin, Y. Chen, S. Hoogland, A.

Rothenberger, K. Katsiev, Y. Losovyj, X. Zhang, P. A. Dowben, O. F. Mohammed, E. H.

Sargent, O. M. Bakr, Science 2015, 347, 519.

[46] Y. Liu, Z. Yang, D. Cui, X. Ren, J. Sun, X. Liu, J. Zhang, Q. Wei, H. Fan, F. Yu, X.

Zhang, C. Zhao, S. F. Liu, Adv. Mater. 2015, 27, 5176.

[47] Z. B. Wang, M. G. Helander, M. T. Greiner, J. Qiu, Z. H. Lu, Phys. Rev. B 2009, 80,

235325.

[48] W. Rehman, R. L. Milot, G. E. Eperon, C. Wehrenfennig, J. L. Boland, H. J. Snaith, M. B.

Johnston, L. M. Herz, Adv. Mater. 2015, 27, 7938.

[49] Q. Wang, Q. Dong, T. Li, A. Gruverman, J. Huang, Adv. Mater. 2016, 28, 6734.

[50] Q. Wang, C. Bi, J. Huang, Nano Energy 2015, 15, 275.

25

Table of Contents:

A dual-source precursor approach is developed to fabricate the high-quality and mirror-like

mixed-cation perovskite without involving additional anti-solvent process. By integrating the

perovskite films into the planar heterojunction solar cells, a power conversion efficiency of

20.15% is achieved with negligible hysteresis effect. A stabilized power output approaching 20%

is obtained at the maximum power point.

Keywords: Dual-Source Precursor Approach, Perovskite Solar Cells, Inverted Planar Heterojunction, Highly Efficient, Purely Stable Mixed-Cation Perovskite

Deying Luo, Lichen Zhao, Jiang Wu, Qin Hu, Yifei Zhang, Zhaojian Xu, Yi Liu, Tanghao Liu, Ke Chen, Wenqiang Yang, Wei Zhang, Rui Zhu* and Qihuang Gong

Dual-Source Precursor Approach for Highly Efficient Inverted Planar Heterojunction Perovskite Solar Cells

26